Process for case hardened dies in improving die life

ABSTRACT

Steel alloys susceptible to case and core hardening comprise 0.05 to 0.24 wt % carbon; 15 to 28 wt % cobalt and 1.5 to 9.5 wt % in nickel, small percentages of one or more additives: chromium, molybdenum, and vanadium; and the balance iron. Carburized roll form and punch dies and tools made from case hardened steel alloys with a reduced hardness core provide high wear and fatigue resistance as well as improved contact and bending fatigue resistance thereby avoiding premature failure and extending the useful life of such tools.

This patent application is a divisional of U.S. patent application Ser. No. 09/722,769, filed Nov. 27, 2000 now U.S. Pat. No. 6,379,475, for which priority is claimed. U.S. patent application Ser. No. 09/722,769 is a continuation-in-part of application Ser. No. 09/664,021 filed Sep. 19, 2000, now U.S. Pat. No. 6,491,767 entitled “Case Hardened Dies For Improved Die Life” which is a continuation-in-part of application Ser. No. 09/239,131 filed Jan. 28, 1999, now U.S. Pat. No. 6,176,946 entitled “Advanced Case Carburizing Secondary Hardening Steels” which is a continuation of Provisional Application No. 60/072,834 filed Jan. 28, 1998 entitled “Advanced Case Carburizing Secondary Hardening Steels”, and for which priority is claimed with respect to said applications.

ACKNOWLEDGMENT OF FUNDING

The original subject matter of the parent application was funded, at least in part, by the Army Research Office, Grant No. DAAH04-96-1-0266. The subsequent continuation-in-part subject matter herein was not funded by Grant No. DAAH04-96 -1-0266

FIELD OF INVENTION

This invention relates to a new class of steel alloys especially useful for the manufacture of case hardened steel gears and other products, in particular, blades such as skating blades made from case hardened steel alloys.

BACKGROUND OF THE INVENTION

This invention relates to a new class of steel alloys especially useful for the manufacture of case hardened steel gears and other products, in particular, blades such as skating blades made from case hardened steel alloys.

Currently, there are a number of high performance steels on the market. Many of these materials utilize primary carbides to achieve high surface hardness and others use stage one or stage three tempered conditions with epsilon carbide or cementite strengthening. Primary carbides are formed when the carbon content exceeds the solubility limit during the solution treatment, and large alloy carbides precipitate. This is the case for secondary hardening steels using alloy carbide strengthening for greater thermal stability to improve properties such as scoring resistance.

However, primary carbide formation can have a detrimental impact on both bending and contact fatigue resistance. Formation of primary carbides can also make process control difficult for avoidance of undesirable carbide distributions such as networks. In addition, primary carbide formation in current steels can lead to a reversal in the beneficial residual compressive stresses at the surface. This is due to a reversal of the spatial distribution of the martensite start temperature due to the consumption of austenite stabilizing elements by the primary carbides. In applications of sliding wear, the formation of primary carbides may be beneficial.

Thus, there has developed a need for case hardenable steel alloys which do not rely upon primary carbide formation, but which provide secondary hardening behavior for superior thermal stability. This invention provides a new class of steel alloys meeting this requirement, while exploiting more efficient secondary hardening behavior to allow higher surface hardness levels for even greater improvements in fatigue and wear resistance. Rolls for manufacturing processes utilizing such steels are projected to have more uniform and enhanced performance characteristics derived through simplified manufacturing technologies and to also have performance characteristics which are more predictable and reproducible and lead to longer life of the tool.

SUMMARY OF THE INVENTION

Briefly, the present invention comprises a class of case hardenable steel alloys in the form of roll form dies with carbon content in the range of about 0.05 wt % to about 0.24 wt % in combination with a mixture of about 15 to 28 wt % cobalt, 1.5 to 9.5 weight 5% nickel, 3.5 to 9.0 wt % chromium, up to 3.5 wt % molybdenum, and up to 0.2 weight percent vanadium.

The microstructural features are a Ni—Co lath martensite matrix steel strengthened by M₂C carbides typically containing Cr, Mo and V. Typical processing of this class of steels includes case carburizing, solution treatment, quenching, and tempering, although due to the high alloy content, quenching may not-be required. Case carburizing produces a gradient in the volume fraction of the M₂C carbides and results in a concomitant increase in hardness and promotes a surface residual compressive stress. The efficiency of the M₂C strengthening response allows this class of steels to achieve very high surface hardnesses with limited soluble carbon content. Thus, this class of steels have the ability to achieve very high surface hardnesses without the formation of primary carbides.

Typical advantages of this class of alloys include ultrahigh case hardness leading to superior wear and fatigue resistance, superior core strength and toughness properties, optional air hardening resulting in less distortion, and higher thermal resistance.

This new class of secondary hardening steels are matrix steels utilizing an efficient M₂C precipitate strengthening dispersion. Because of the efficiency of this strengthening dispersion, a superior combination of properties can be attained for many applications on a situation-by-situation and product-by-product basis. For example, in situations where the desired surface properties are similar to current materials, the core strength and toughness can be superior. In applications where superior surface properties are desired, the disclosed steels can easily outperform typical materials while maintaining normal core properties, and in applications which require corrosion resistance, these new steels can provide stainless properties with surface mechanical properties similar to typical non-stainless grades.

Thus, an object of the invention is to incorporate desirable properties resulting from the class of alloys disclosed in various products.

A further object of the invention is to provide roll form dies, die punches, stamping dies, or draw dies made from case hardened steel alloy materials wherein the surface hardness of the dies surface and the core hardness of the dies are controlled to maximize performance and to provide uniform and reproducible characteristics.

Another object is to provide die forms for rolling, punching, stamping or drawing made from steel alloy materials disclosed wherein the flexibility, hardness, sharpness of the die and other die characteristics are controllable and reproducible.

These and other objects, advantages and features of the invention will be set forth in the detailed description which follows.

BRIEF DESCRIPTION OF THE DRAWING

In the detailed description which follows, reference will be made to the drawing comprised of the following figures:

FIG. 1 is a graph correlating hardness to precipitation driving force for experimental and predicted results;

FIG. 2 is a graph correlating precipitation half completion time and half completion-coarsening rate constant for experimental and predicted results;

FIG. 3 is a graph which correlates calculated segregation free energy difference with experimental embrittlement potency;

FIG. 4 is a flow block diagram of the total system structure of the alloys of the invention;

FIG. 5 is a graph depicting the relationship between cobalt and nickel content for a 200° C. Ms temperature for the alloys of the invention;

FIG. 6 is a pseudo-ternary diagram as a function of chromium, molybdenum and vanadium at 0.55 wt % carbon with regard to alloys of the invention at 1000° C.; and

FIG. 7 is a graph comparing hardness of the steel alloys of the invention with conventional carburized alloys.

FIG. 8 is a graph containing Falex wear test data for steel alloys of the invention in comparison with conventional 8620 steel.

FIG. 9 is a graph of NTN 3 ball-on-rod rolling contact fatigue data for alloys of the invention in comparison with conventional M50 bearing steel;

FIG. 10 is an isometric view of a typical roll form die fabricated from a carburized case hardened steel alloy; and

FIG. 11 is a cross-sectional view of the die of FIG. 10 taken along the line 11—11.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

The steel alloys of the invention were developed using various modeling techniques followed by experimental confirmation or testing. An important component of the modeling is the application of a thermochemical data bank and software system. The system or program employed uses 5thermodynamic assessments from binary, ternary, and quaternary systems to extrapolate to higher order multicomponent systems. Equilibria, constrained equilibria, and driving forces can be calculated as functions of composition, chemical potential, as well as other user defined functions. To apply this information to the modeling of highly nonequilibrium processes of interest in real alloys, the dynamic nature of phase transformations in terms of thermodynamic scaling factors are described and then evaluated by the thermochemical software. Thus, hypothetical steel compositions were the subject of an initial computational model involving the precipitation of M₂C carbides—leading to a secondary hardening response in ultrahigh-strength steels. A second effort employed a published thermodynamics-based model for the non-linear composition dependence of the martensite start temperature. A third modeling effort involves the application of quantum 15 mechanical calculations to the production of hypothetical compositions with the goal of achieving improved resistance to hydrogen embrittlement and intergranular fracture. Modeling techniques were then followed by testing of the optimized alloys. Following is a discussion of modeling technique considerations.

Secondary Hardening

Ultrahigh-strength (UHS) secondary hardening steels are strengthened by the precipitation of coherent M₂C carbides during tempering. In high Co steels in which dislocation recovery is retarded, the M₂C carbides precipitate coherently on dislocations and provide the characteristic secondary hardening peak during tempering. A wide range of techniques are utilized to gather experimental information across a complete range of size and time scales of interest. Atom-probe field-ion microscopy (APFIM), transmission electron microscopy (TEM), small angle neutron scattering (SANS), and X-ray diffraction (XRD) techniques provide information on the size, shape, composition, and lattice parameters of the M₂C precipitates as well as the resulting hardness values spanning tempering times of less than an hour to more than a thousand hours. This study identified that the precipitation was well described by a theory developed by Langer and Schwartz for precipitation at high super saturation in which the growth regime is suppressed and precipitation occurs by a process of nucleation and coarsening, maintaining a particle size close to the critical size.

Based on these investigations, two important scaling factors are identified. The initial critical 15 nucleus determines the size scale of the precipitates throughout the precipitation reaction and the coarsening rate constant determines the precipitation time scale. The peak hardness in an ultra high strength steel commonly occurs at the particle size corresponding to the transition from particle shearing to Orowan bypass. It is also advantageous to bring the M₂C precipitation to completion in order to dissolve all of the transient cementite which otherwise limits toughness and fatigue life.

Therefore, the smaller the initial critical particle size, the closer completion of precipitation occurs to peak hardness and more efficient strengthening is obtained. The time scale of precipitation is also important due to the kinetic competition between the secondary hardening reaction and the segregation of impurities to the prior austenite grain boundaries leading to intergranular embrittlement.

The initial critical nucleus size scales inversely with the thermodynamic driving force for precipitation. In the case of the M₂C carbide it is important to include the influence of prior cementite formation and coherency on this quantity. The coherency elastic self energy can be evaluated by the calculation of an anisotropic ellipsoidal inclusion using the equivalent Eigenstrain method and the impact of solute redistribution on the resulting stress distribution is addressed by using open-system elastic constants. By relating the coherency strain to composition via the compositional dependence of the particle and matrix lattice parameters, the composition dependence of the elastic self energy is determined in a form compatible with the thermodynamic software. The linear elastic self energy calculation represents an upper limit and a correction factor is used to fit the precipitation composition trajectories of a large set of experimental alloys.

The impact of prior cementite precipitation is accounted for by the calculation of the coherent driving force. in the presence of the carbon potential due to para-equilibrium cementite. This para-equilibrium carbon potential is defined by an equilibrium between the matrix and cementite in which the substitutional species are held constant and only the interstitial carbon is allowed to partition.

In this approximation the cementite acts as a carbon source at constant chemical potential.

FIG. 1 represents the level of agreement of the strengthening response of the model alloys with the above model. The model alloys contain 16 wt % Co, 5 wt % Ni, and 0.24 wt % C with varying amounts of the carbide formers Cr, Mo, and, in a few cases, V. The nickel content is chosen to eliminate austenite precipitation during tempering which otherwise complicates the hardening response. In FIG. 1, the peak hardness during tempering at 510° is plotted against the driving force for precipitation of the coherent M₂C carbide in the presence of para-equilibrium cementite.

The open circles represent alloys containing V. The relationship demonstrates the ability to predict peak hardness values within approximately +/−25 VHN in this class of alloys.

The time scale of precipitation at high super saturations, according to the Langer-Schwartz treatment, scales with the coarsening rate of the particle distribution. The modeling pursued in this work expands upon the Lifshitz-Slyozov and Wager (LSW) theory, describing the coarsening of spherical particles in a binary system, with the intent of removing the binary restrictions of the LSW theory and reformulating it in a manner compatible with the multicomponent thermodynamics of the software and data bank system.

The result of this analysis characterizes the coarsening rate of a particle of average size as a function of the multicomponent diffusion coefficients, the equilibrium partitioning coefficients, and the second derivatives of the Gibbs free-energy evaluated at the equilibrium state. The surface energy and molar volume are taken to be composition independent and are considered constant. In this form, the coarsening rate constant is the result of an asymptotic analysis and is only representative at very long time scales and very close to equilibrium. This is certainly not the case for the precipitation of the M₂C carbide at high super saturation. The matrix content of alloy is far from equilibrium during much of the precipitation process, approaching equilibrium only near completion. This effect is more severe for alloys containing stoichiometric quantities of carbide formers as measured by the relative difference in the matrix alloy content during precipitation and at equilibrium.

During precipitation in a stoichiometric alloy, the alloy matrix content is of the same order as the overall alloy content, while at equilibrium, the matrix alloy content is very small. To define a coarsening rate constant more representative of the conditions present during the precipitation process, a coarsening rate is evaluated at the point when the volume fraction of the precipitate is one-half of the equilibrium value. This is achieved by calculating the coherent equilibrium for the M₂C carbide, and then, adding energy to the M₂C phase to account for capillarity until the amount of the phase is half of the equilibrium value.

The coarsening rate is then calculated from the thermodynamic properties of this state. FIG. 2 represents the correlation between the precipitation half-completion time and the half-completion coarsening rate constant of the model alloys for which this data is available.

Ms Temperatures

To predictively control the spatial distribution of martensite-start (Ms) temperatures in the carburized steels to achieve fully martensitic structures with controlled residual stress distributions, a published model was employed. The thermodynamics-based nucleation-kinetic model was calibrated to the composition-dependence of measured Ms temperatures using both literature data and assessments of experimental multi component alloys.

Interfacial Cohesion

Intergranular embrittlement phenomena such as hydrogen embrittlement are undesirable in the intended alloys. Embrittlement of ultrahigh-strength steels is associated with the prior segregation to the grain boundary of impurities such as P and S. A thermodynamic treatment of this phenomenon by Rice and Wang illustrates that the potency of a segregating solute in reducing the work required for brittle fracture along a boundary is linearly related to the difference in the segregation energy for the solute at the boundary and at the free surface. Specifically, a solute with a higher segregation energy at the free surface will be an embrittler while a solute with a higher segregation energy at the grain boundary will enhance intergranular cohesion. A survey of reported segregation energies and embrittling potency (reported as the shift in the ductile-to-brittle transition temperature per atomic percent solute on the grain boundary) in Fe-base alloys demonstrates these general trends; however, the experimental difficulty of surface thermodynamic measurements gives ambiguous values for some solutes.

First principle calculations were used to determine the total energy of atomic cells representing the Fe Σ3 [110](111) grain boundary and (111) free surface with a monolayer of an impurity solute present. The calculations were accomplished with the full-potential linearized plane wave (FLAPW) total energy technique. The atomic structure in each case was relaxed to find the minimum energy state. The results of these calculations include not only the segregation energies responsible for the embrittling or cohesion enhancing effects of segregating solutes, but the underlying electronic structure of the solutes in the boundary and surface environments. A comparison of the directional covalent electronic structure between B, a strong cohesion enhancer, and P, a strong embrittler, indicates the strong bonding of the B atom across the boundary plane associated with hybridization of the B 2 p electrons with the Fe d band. This directional bonding is not seen in the case of the P atom which does not significantly hybridize with Fe.

The results of the first principle calculations were correlated to the experimental embrittling potency in FIG. 3. The difference between grain boundary and free surface segregation energies, calculated by electronic structure calculations, and the experimentally observed shift in the ductile-to-brittle-transition-temperature are plotted for C, B, P and S solutes. The C and B are shown as cohesion enhancers, P and S as embrittlers. The computed energy differences are in excellent agreement with the observed effects on interfacial cohesion.

MATERIALS DESIGN

Background

Design considerations for high performance rolls, punches, stamping and draw dies for various applications include the desire to provide very high surface hardness while maintaining material ductility for shock and flow tolerance. Critically, the portion of the tool (die) which contacts the work surface must be very hard, i.e., greater than 58 Rockwell C hardness. Heretofore, monolithic tool steels with a hardness ranging from about 58 to 62 Rockwell C hardness have been used. However, such steels are brittle and do not exhibit a long term work life span. Alternatively, a reduced hardness alloy may be coated with wear resistant or hard coatings to improve tool life and surface lubricity. This involves potentially expensive coating techniques. Alternatively, very high hardness ceramic or powder metal dies can be used, but they are very expensive and have limited ductility making them hard to manufacture and prone to damage if not handled properly. The alternative of the present invention provides benefits not available from these prior techniques. It involves case hardening by carburizing or nitriding an alloy taken from a special group of alloys in the form of the roll form die, punch, stamping or draw die wherein the form is hardened greater than about 58 Rockwell C with the core maintained at a lesser hardness and wherein the surface is hardened at a depth of at least about 1 millimeter in the region of the die wear surface. The core is typically less than about 53 Rockwell C.

Punch dies or stamping dies may also be made in accord with the described procedure and formulation. The work surface will again be case hardened by a carburizing, or nitriding process, or both, typically of the select steel disclosed thereby providing a gradient between a 1 to 3 millimeter layer of high hardness (>62 Rockwell C) work surface and a relatively softer core or tool matrix (typically <53 Rockwell C).

Analysis

The systems analysis of the case-core secondary hardening steel system is the first step in the design process. FIG. 4 illustrates the total processing/structure/properties/performance system structure for high power-density gears manufactured by three alternative processing routes, conventional forged ingot processing, near net shape casting and powder processing. Case hardenable secondary hardening gear steels are a subsystem of this flow-block diagram and are the focus of this disclosure. The sequential processing steps dictate the evolution of the case and core microstructures which determine the combination of properties required for the overall performance of the system. Both the case and core consist of microstructures of a martensite with high Co, for dislocation recovery resistance essential for efficient secondary hardening, and Ni, for cleavage resistance. Strengthening is provided by the coherent precipitation of fine M₂C carbides on dislocations. This secondary hardening reaction dissolves the transient cementite and it brings the precipitation reaction to completion in order to eliminate cementite for high toughness and fatigue resistance. The grain refining dispersion has a double impact on toughness. By limiting grain growth at high temperature during solution treatment, brittle intergranular modes of fracture are inhibited.

The grain refining particles also play an important role in the ductile microvoid nucleation and coalescence fracture behavior. Thus, it is desired to preserve adequate volume fraction and size to pin the grain boundaries while choosing the phases with greatest interfacial cohesion. Also desirable is the control of the grain boundary chemistry to avoid intergranular embrittlement (such as by hydrogen embrittlement) in association with prior segregation of embrittling impurities.

During tempering, impurities segregate to the grain boundaries and in the case of P and S reduce the interfacial cohesion of the boundary promoting intergranular embrittlement. A number of methods are used to avoid this problem. Gettering compounds can be utilized to tie up the impurities in stable compounds reducing the segregation to the grain boundary. In order to produce the most stable compounds, however, rapid solidification processing is required. Within their solubility limits, additional segregants such as B can be deliberately added to enhance intergranular cohesion, and the precipitation rate for the secondary hardening reaction can be increased to limit the time at temperature for harmful grain boundary segregation.

Design

As a first design step, core and case carbon levels required for the desired hardness are estimated. This is done by fitting data for existing secondary hardening Ni—Co steels to an Orowan strengthening model and extrapolating to the desired strength. It is estimated that a core carbon content of 0.25 wt % and a case carbon content of 0.55 wt % is needed to provide the desired core and case hardness in this Ni—Co steel.

The next step is to determine the matrix composition of the Fe, Ni, and Co. In order to produce the desired lath martensite morphology, an Ms temperature of 200° C. or above is required. Using the nucleation kinetic model for the compositional dependence of the Ms temperature, the variation with Ni and Co content is determined. This result is illustrated in FIG. 5 for the case carbon content using a preliminary composition of carbon formers equal to 5 wt % Cr, 0.5 wt % Mo and 0.0 wt % V. Since the case has a higher carbon content than the core, the core will possess a higher Ms temperature than the case. In FIG. 5 the Co and Ni content required to fix the _(s)M temperature at 200° C. is indicated. Since a high Ni content is desired to avoid cleavage fracture, a composition containing 25 wt % Co was chosen. This allows the highest possible Ni content, approximately 3.5 wt %, to be used. These calculations are later repeated for consistency when the composition of the carbide formers is further refined.

To define the optimal composition of the carbide formers a number of design constraints are applied. The total amount of carbide formers in the alloy must be greater than that required to consume the carbon present in the case. This lower limit insures that, at completion, embrittling cementite is completely converted to M₂C carbide. In order to reduce grain boundary segregation, the precipitation rate is maximized. This allows the shortest possible tempering time. The coherent precipitation driving force is maximized to provide a small critical particle size for the M₂C and more efficient strengthening. Finally, the solution temperature is limited to 1000° C. This allows Cr, Mo and V containing carbides such as M₂₃C₆, M₇C₃, MC and M₆C to be dissolved at reasonable processing temperatures while maintaining very fine scale TiC carbides to act as the grain refining dispersion.

Calculations for the precipitation rate constant indicate low Mo compositions are favorable, while driving force calculations have demonstrated the highly beneficial effect of higher V contents. The solubility constraints are presented by the diagram in FIG. 6. Here the equilibrium phase fields at 1000° C. are given as a function of Cr and V content. The Mo content is determined by the stoichiometry requirements, the matrix composition is taken from the earlier calculations, and the carbon content represents the case composition. The point on the diagram within the single phase FCC field that maximizes the V content and minimizes the Mo content represents the compromise fulfilling the design criteria. This composition is 4.8 wt % Cr, 0.03 wt % Mo, and 0.06 wt % V. A recalculation of the matrix composition using the final carbide formers results in an alloy composition of Fe-25 Co-3.8 Ni-4.8 Cr-0.03 Mo-0.06 V-0.55 (case)/0.25 (core) C (in wt %).

Consistent with the model predictions of FIG. 3, a soluble boron addition of 15-20 ppm is added to enhance intergranular embrittlement resistance.

EXAMPLES

A 17 lb. vacuum induction heat of the above composition was prepared from high purity materials. The ingot was forged at 1150° C. in a bar 1.25″ square by 38″ long. The M_(s) temperature of the alloy was determined from dilatometery and found to agree with model predictions. The solution treatment response of the alloy was determined from hardness measurements in the stage I tempered condition. The optimum processing conditions for the core material was determined to be a 1050° C. 1 hour solution treatment followed by an oil quench and a liquid nitrogen deep freeze. After optimal solution treatment, a 12 hour temper at 482° C. results in the desired overaged hardness of 55 R_(c) for the core material. The material was then plasma carburized and processed using these parameters. The C potential, temperature and time used in the carburizing treatment were determined from simulations with multicomponent diffusion software to provide the target surface carbon content of 0.55 wt % and a 1 mm case depth. The curve labeled C2 in FIG. 7 represents the hardness profile achieved for the carburized sample. A surface hardness of 67 HR_(c) and a case depth of 1 mm are obtained.

Using techniques and processes of this nature, the following alloys set forth in Table 1 were developed and tested:

TABLE 1 Alloy Fe Co Ni Cr Mo V C (Core) A1 Bal. 18 9.5 3.5 1.1  0.08 0.20  C2 Bal. 25 3.8 4.8 0.03 0.06 0.237 C3 Bal. 28 3.25-3.15 5.0 1.75-2.50 0.025 0.05-0.18 CS1 Bal. 15 1.5 9.0 0.0  0.2  0.05-0.20

The first, AI, is targeted as a replacement for current gear materials in applications where component redesign is not feasible but higher core strength and toughness is needed. As such, A1 has surface wear properties similar to current commercial properties, but possesses superior core toughness and strength 54 HRC and a K_(Ic) of>75 Ksivin. The second alloy C2 corresponds to the prototype alloy just described. The third alloy, C3, pushes the surface properties s far as possible while maintaining adequate core strength and toughness. As also shown by the hardness profiles of FIG. 7, the alloy has reached a surface hardness corresponding to HRC 69. Wear tests for the carburized material in a standard Falex gear simulator show much reduced weight loss compared to standard carburized 8620 steel in FIG. 8. A ball-on-rod rolling contact fatigue test (NTN type) conducted at 786 ksi Hertzian contact stress indicates an order of magnitude increase in LIO fatigue life compared to M50 bearing steel as shown in FIG. 9. The fourth alloy in Table 1, CS1, represents a stainless variant of this class of alloy. Targeted to match the surface properties of standard non-stainless gear and bearing materials with sufficient core strength and toughness, the alloy has achieved corrosion resistance better than 440C by anodic polarization conducted in distilled water with a neutral ph. (sucrose added for electrical conductivity). Similar relative behaviors were demonstrated in 3.5% NaCl solution. In salt fog tests, CS1 outperformed 440C and commercial carburizing stainless steels, the performance gap widened when the tests were completed on samples in the. carburized condition. The carburized alloy achieved surface mechanical properties equivalent to A1 while maintaining corrosion resistance. In RFC tests of the type represented in FIG. 9, both A1 and CS1 showed L₁₀ fatigue life equal or superior to the M50 bearing steel. The table (Table 2) below summarizes the performance achieved in the four alloys:

TABLE 2 Rolling Core Core Surface Bending Contact Corrosion Alloy Hardness¹ Toughness² Hardness³ Fatigue⁴ Fatigue⁵ L₁₀ Resistance⁶ A1   54 R_(c) >75 ksiin >61 R_(c)   ≧EN36C ≧M50 NA C2 ≦58 R_(c) adjustable 67 R_(c) NA In process NA C3 ≦59 R_(c) adjustable 69 R_(c) NA 10 × M50 NA CS1 ≦53 R_(c) ≧25 ksiin 63 R_(c) NA ≧M50 >440C adjustable ¹Hardness determined by ASTM E18. ²Core toughness determined by ASTM E813. ³Surface hardness determined by E384. ⁴Bending fatigue determined using 4-point bend testing. ⁵Rolling contact fatigue determined by NTN 3 ball-on-rod techniques. ⁶Corrosion resistance determined by anodic polizations and salt for testing.

Each alloy has a surface hardness exceeding and core hardness exceeding prior art compositions achieved at a lower cost. Variants of the disclosed alloys are set forth in Table 3 (Nominal Compositions in wt. %):

TABLE 3 Alloy Fe Co Ni Cr Mo V C(ore) A1 (modified) Bal 18 9.5 3.5 1.1 0.08 0.16 C3 (modified) Bal 27.8-28.2 2.9-3.1 5.0-5.2 2.4-2.6 0.18-.025 0.07 CSI (modified) Bal 15 1.5 9.0 0.2  0.08

The variants are considered equivalent to the unmodified alloys of Table 1 and are within the class of alloys comprising the alloys of the invention.

It has also been discovered that the disclosed alloys may be subjected to a nitriding process as well as a carburizing process and other case hardening processes such as induction heating to enhance surface hardness. For example the alloy C3 set forth above may be nitrided during the tempering step to increase surface hardness by 10% or more.

FIGS. 10 and 11 depict the incorporation of steels of the type discussed above in a die.

Various types of dies may be fabricating using the described steels, including roll form dies as well as punch dies. Sheet forming dies or stamping dies may also be made using the described steels.

Referring to FIG. 10, there is illustrated a representation of a roll form die 48. Draw dies used to form wire products may also be made using the described steels. Die 48 typically includes a shaped form or surface 50 which is engaged against a material (not shown) which is to be formed or shaped by the die 48. Typically, the die 48 is incorporated in or defines a roll and rotates as depicted by the arrow in FIG. 10 against the material to shape and form the material. The surface 50 desirably has a hardness greater than 58 Rockwell C and a case depth of preferably more than about 0.1 millimeter. More preferably, the surface hardness exceeds 65 Rockwell C and the depth of the case of the hardened die is greater than about 1 millimeter and preferably in the range of about 1 to 3 millimeters. The surface may be hardened by carburizing, nitriding, induction heating or other methods to improve the strength of the material at the surface and thereby achieve case hardening.

Importantly, the core of the die has a lower hardness in the range of less than about 53 Rockwell C. This improves the wear and fatigue resistance at the surface while maintaining a lower hardness core which is ductile. Thus, there is a gradient in hardness with respect to the surface and core with the high hardness surface providing high wear and fatigue resistance while the ductile core provides shock and flaw tolerance. Also, the compressive residual stresses in the case improve the contact and bending fatigue resistance at the surface which can otherwise lead to premature failure of the tool.

FIG. 11 illustrates in cross section, a typical construction of such a die 48 wherein the surface 50 comprises the hardened portion of the die 48 inasmuch as the curved or shaped surface 50 engages the material which is being shaped by the die. Besides carburizing, nitriding, induction hardening, other surface modification techniques may be used followed by heat treatment to produce the desired case and core properties. The result is a very high hardness for the surface, particularly in high wear areas as well as beneficial residual compressive stresses. The strength of the core can also be controlled to provide sufficient strength to withstand the internal body stresses applied to the tool. The lower hardness of the core also provides improved flaw tolerance and ductility. As a result, the tool has significantly longer useful life.

Of course, the drawing depicts a roll form die but other types of dies including punch dies, sheet fabrication dies, draw dies and the like may be fabricated. It is noted that with roll form dies, the resultant die and the process of using the die provides improved wear, notch bending, fatigue resistance and contact fatigue resistance as well as shock and flaw tolerance. Similar benefits are observed with sheet fabrication and punch dies. The dies exhibit a longer useful life and can be used to produce more complex shapes as well as punch or form dies with higher aspect ratio features.

These benefits provide that a single die may be utilized to provide more complex shapes and cuts.

The manufacture of the dies of the invention are easier because the machining of the shape of the tool can be completed in the soft state and then case hardened to high hardness, considering the distortion involved in such a process.

Various aspects of the invention may therefore be altered without changing the form and scope of the invention. Thus, the invention is to be limited only by the following claims and equivalents thereof. 

What is claimed is:
 1. A process for forming materials comprising the steps of: providing a tool from the group consisting of dies and punches including a core and a work surface made from a steel alloy comprised of about 0.008 to 0.24% by weight carbon, about 15 to 28% by weight cobalt, about 1.5 to 9.5% by weight nickel, and one or more additives taken from the group consisting of chromium, molybdenum and vanadium, and the balance iron; case hardening the work surface of the tool to a depth of at least 0.1 millimeters at a hardness of at least 65 Rockwell C; and working the material with the tool by engaging the material with the work surface.
 2. The process of claim 1, wherein the depth is about 1 to 3 millimeters.
 3. A process for forming materials comprising the steps of: providing a tool from the group consisting of dies and punches including a core and a work surface made from a steel alloy comprised of about 0.008 to 0.24% by weight carbon, about 15 to 28% by weight cobalt, about 1.5 to 9.5% by weight nickel, and one or more additives taken from the group consisting of chromium, molybdenum and vanadium, and the balance iron; case hardening the work surface of the tool to a depth of at least 0.1 millimeters at a hardness in the range of about 58 to 62 Rockwell C; and working the material with the tool by engaging the material with the work surface.
 4. The process of claim 3, wherein the depth is about 1 to 3 millimeters. 